High gamma prime nickel based superalloy, its use, and method of manufacturing of turbine engine components

ABSTRACT

The specification relates to a high gamma prim nickel based superalloy, its use and a method of manufacturing of turbine engine components by welding, 3D additive manufacturing, casting and hot forming, and the superalloy comprises by wt %: from 9.0 to 10.5% Cr, from 16 to 22% Co, from 1.0 to 1.4% Mo, from 5.0 to 5.8% W, from 2.0 to 6.0% Ta, from 1.0 to 4.0% Nb provided that total content of Ta and Nb remains with a range from 3.0 to 7.0%, from 3.0 to 6.5% Al, from 0.2 to 1.5% Hf, from 0.01 to 0.2% C, from 0 to 1.0% Ge, from 0 to 1.0 wt. % Si, from 0 to 0.2 wt. % Y, from 0 to 0.015 wt. % B, from 1.5 to 3.5 wt. % Re, and nickel with impurities to balance.

CROSS-REFERENCE TO RELATED APPLICATION

This application claims the benefit of and priority to Chinese PatentApplication No. 201911060106.X filed Nov. 1, 2019 under the title HIGHGAMMA PRIME NICKEL BASED SUPERALLOY, ITS USE, AND METHOD OFMANUFACTURING OF TURBINE ENGINE COMPONENTS. The content of the abovepatent application is hereby expressly incorporated by reference intothe detailed description hereof.

FIELD

The present disclosure relates to high gamma prime (γ′) nickel-basedsuperalloy can be used for laser beam (LBW), plasma (PAW), micro-plasma(MPW), gas tungsten arc welding (GTAW), electron beam (EBW) welding and3D additive manufacturing (AM), as well as for manufacturing of turbineengine components and other articles by casting and hot forming.

BACKGROUND

Most turbine blades of aero and industrial turbine engines aremanufactured from nickel based high gamma-prime (γ′) superalloys thathave unique combination of oxidation and creep properties. However,despite remarkable properties of high γ′ superalloys, engine componentsfrequently require various weld repairs due to creep andthermo-mechanical fatigue cracking, oxidation and hot corrosion damageoccurring during operation of turbine engines. Dissimilar cobalt basedMerl 72 (M72), nickel based René 142 (R142) and René 80 (R80) weldingmaterials have been used for a repair of high (HPT) and low (LPT)pressure turbine blades from 1980-s, refer to A. Gontcharov et al,GT2018-75862, “Advanced Welding Materials and Technologies for Repair ofTurbine Engine Components manufactured of High Gamma Prime Nickel BasedSuperalloys”, Proceedings of ASME Turbo Expo 2018: Turbine TechnicalConference and Exposition, GT2018, Jun. 11-15, 2018, Oslo, Norway(further GT2018-75862) (incorporated herein by reference).

Cobalt based M72 has excellent weldability, ductility and oxidationresistance but low creep properties at temperatures ≥1800° F. as shownin GT2018-75862 and Example 1, which resulted in a premature HPT bladesfailure and unscheduled engine removals. Low creep properties aretypical for most cobalt based alloys and nickel-based superalloys withhigh cobalt content. On the other hand, high γ′ nickel based R142welding wire, which comprises 6.8 wt. % Cr-12 wt. % Co-1.5 wt. % Mo-4.9wt. % W-6.4 wt. % Ta-6.1 wt. % Al-1.5 wt. % Hf-2.8 wt. % Re, that wasdisclosed by Earl W. Ross and Kevin S. O'Hara “Rene 142: High Strength,Oxidation Resistance DS Turbine Airfoil Alloy”, Superalloys 1992, pp.257-265 and created based on the high gamma prime nickel basedsuperalloy as per U.S. Pat. No. 4,169,742 that comprised of: 10-13 wt. %Co, 3-10 wt. % Cr, 0.5-2 wt. % Mo, 3-7 wt. % W, 0.5-10 wt. % Re, 5-6 wt.% Al, 5-7 wt. % Ta, 0.5-2 wt. % Hf, 0.01-0.15 wt. % C, 0.005-0.05 wt. %B, 0-0.1 wt. % Zr with nickel to balance, has excellent creepproperties, but extremely poor weldability. Limited weld repairs ofturbine engine components with R142 have been done only with thepreheating of engine components to high temperature as it wasdemonstrated by Dikran A. Barhanko et al, “Development of Blade TipRepair for SGT-700 Turbine Blade Stage 1, With Oxidation Resistant WeldAlloy”, Proceedings of ASME Turbo Expo 2018, Turbomachinery TechnicalConference and Exposition, GT2018, Jun. 11-15, 2018, Oslo, Norway andAlexandre Gontcharov et al in the previously quoted GT2018-75862article. However, even with the preheating, R142 welds demonstrated poorductility and high propensity to microcracking such that it is unable touse R142 for 3D additive manufacturing.

Nickel based superalloy R80 with the chemical composition as per U.S.Pat. No. 3,615,376, which comprises Ni-15% Cr-9.5% Co-5% Ti-4% W-4%Mo-3% Al-0.17% C, has better weldability but poor oxidation resistanceand can't substitute R142 and M72.

Nickel based superalloys disclosed in CN 105492639, CA 28004402, U.S.Pat. Nos. 4,288,247, 7,014,723, 8,992,669, and 8,992,700 with elevatedto 20-30% Co content can't substitute the high gamma prime R142superalloy as well due to insufficient mechanical properties at ≥1800°F. despite potentially better weldability.

Therefore, there are substantial needs in the development of new highoxidation resistance, high strength and ductility high gamma primenickel-based superalloys that can help to produce crack free welds onsingle crystal (SX) materials at an ambient temperature for repair and3D AM of turbine engine components.

BRIEF DESCRIPTION OF THE INVENTION

We have found that the high gamma prime nickel-based superalloycomprising by wt. % from 9.0 to 10.5% Cr, from 16 to 22% Co, from 1.0 to1.4% Mo, from 5.0 to 5.8% W, from 2.0 to 6.0% Ta, from 1.0 to 4.0% Nbprovided that total content of Ta and Nb remains with a range from 3.0to 7.0%, from 3.0 to 6.5% Al, from 0.2 to 1.5% Hf, from 0.01 to 0.2% C,from 0 to 1.0% Ge, from 0 to 1.0 wt. % Si, from 0 to 0.2 wt. % Y, from 0to 0.015 wt. % B, from 1.5 to 3.5 wt. % Re, and nickel with impuritiesto balance, has excellent weldability at an ambient temperature, goodcombination of mechanical and oxidation properties and can be used forvarious repairs of turbine engine components by a fusion welding and forthe manufacturing of turbine engine components by 3D AM, casting, andhot forming.

In a particular embodiments, the amount of each element can be variedindependently of the other elements. For instance, the high gamma primenickel-based superalloy can contain:

Cr having a range (in wt. %) from 9.3-10.5, 9.6-10.5, 9.9-10.5,9.0-10.2, 9.3-10.2, 9.6-10.2, 9.0-9.9, 9.6-9.9, 9.5-10 or 9.5-10.5, orvariations thereof:

Co having a range (in wt. %) from 16-21; 16-20; 16-19; 17-22; 17-21;17-20; 17-19; 18-22; 18-21; 18-20; or 18-19, or variations thereof:

Mo having a range (in wt. %) from 1.0-1.3; 1.0-1.2; 1.1-1.4; 1.1-1.3;1.1-1.2; 1.2-1.4; or 1.2-1.3, or variations thereof;

W having a range (in wt. %) from 5.1-5.8; 5.2.-5.8; 5.3-5.8, 5.4-5.8,5.0-5.7, 5.1-5.7; 5.2.-5.7; 5.3-5.7, 5.4-5.7, 5.0-5.6, 5.1-5.6;5.2.-5.6; 5.3-5.6, 5.4-5.6, 5.0-5.5, 5.1-5.5; 5.2.-5.5; 5.3-5.5, or5.4-5.5, or variations thereof;

Ta having a range (in wt. %) from 2.5-6.0, 3.0-6.0, 3.5-6.0, 4.0-6.0,2.0-5.5, 2.5-5.5, 3.0-5.5, 3.5-5.5, 4.0-5.5, 2.0-5.0, 2.5-5.0, 3.0-5.0,3.5-5.0, 4.0-5.0, 2.0-4.5, 2.5-4.5, 3.0-4.5, 3.5-4.5, 4.0-4.5, orvariations thereof;

Nb having a range (in wt. %) from 1.5-4.0, 2.0-4.0, 2.5-4.0, 1.0-3.5,1.5-3.5, 2.0-3.5, 2.5-3.5, 1.0-3.0, 1.5-3.0, 2.0-3.0, or 2.5-3.0, orvariations thereof, and wherein the total Ta+Nb content has a range (inwt. %) from 3.5-7.0, 4.0-7.0, 4.5-7.0, 5.0-7.0, 3.0-6.5, 3.5-6.5,4.0-6.5, 4.5-6.5, 5.0-6.5, 3.0-6.0, 3.5-6.0, 4.0-6.0, 4.5-6.0, 5.0-6.0,or 5.5-6.0, or variations thereof:

Al having a range (in wt. %) from 3.5-6.5, 4.0-6.5, 4.5-6.5, 3.0-6.0,3.5-6.0, 4.0-6.0, 4.5-6.0, 3.0-5.5, 3.5-5.5, 4.0-5.5, or 4.5-5.5, orvariations thereof:

Hf having a range (in wt. %) from 0.4-1.5, 0.6-1.5, 0.8-1.5, 0.2-1.2,0.4-1.2, 0.6-1.2, 0.8-1.2, 0.2-1.0, 0.4-1.0, 0.6-10, or 0.8-1.0, orvariations thereof;

Each of Ge and Si (independently) (in wt. %) from trace amount to 1.0,0.2-1.0, 0.4-1.0, 0.6-1.0, trace amount to 0.8, 0.2-0.8, 0.4-0.8,0.6-0.8, trace amount to 0.6, 0.2-0.6, or 0.4-0.6, or variationsthereof;

Y having a range from trace amount to 0.2, 0.05-0.2, 0.1-0.2, traceamount to 0.1, 0.05-0.1, or variations thereof;

B having a range from trace amount to 0.015, 0.005-0.015, 0.010-0.015,0-0.010, trace amount to 0.010, or 0.005-0.010, or variations thereof:

C having a range (in wt. %) from 0.05-0.2, 0.1-0.2, 0.01-0.15,0.05-0.15, or 0.1-0.15, or variations thereof, or

Re having a range (in wt. %) from 2.0-3.5, 2.5-3.5, 1.5-3.0, 2.0-3.0, or2.5-3.0, or variations thereof.

The above variations, or their variations, can be combined in line withdescription of the alloy disclosed herein.

In one embodiment, the high gamma prime nickel-based superalloycomprises total amount of germanium and silicon within the range from0.9 to 1.1 wt. %.

In another embodiment, the current superalloy is selected from amongwelding wire, welding powder, equiaxed or directionally solidifiedturbine engine component, repaired turbine engine component, and articleproduced by hot forming.

In another aspect of the present invention, a method of manufacturing aturbine engine component is provided, wherein it comprises a step ofusing the high gamma prime nickel-based superalloy of the presentinvention.

Herein, “manufacturing a turbine engine component” refers to themanufacturing from the raw material and/or repairing an old turbineengine component such that it can be used as a new one.

Turbine engine components and other articles manufactured from theinvented superalloys with the preferable chemical composition aresubjected to heat treatment, which is different from the heat treatmentof R142 superalloy, which produce the best properties by annealing at2300° F. for 2 hours followed by primary aging at 1975° F. for 4 hoursand secondary aging at 1650° F. for 4 hours, and includes annealingwithin the temperature range from 2190° F. to 2290° F. for 1-2 hours,primary aging within the temperature range from 1975° F. to 2050° F. for2-4 hours, and secondary aging within the temperature range from 1300°F. to 1500° F. for 16-24 hours aiming to maximize mechanical propertiesof the developed superalloy by the aging that results in a precipitationof γ′ phase.

In an embodiment, manufacturing of turbine engine components by castingcomprises an additional step of a hot isostatic pressure treatment of aningot at a temperature of 2200-2290° F., pressure of 15-20 KSI(102.6-136.8 MPa) for 2-6 hours prior to annealing.

Manufacturing of turbine engine components as per another embodimentcomprises at least two consecutive steps of the annealing of the ingotat 2190° F. to 2290° F. for 1-2 hours followed by the hot forming withthe temperature range from 1500° F. to 1800° F. by a plastic deformationby 5-80% and final heat treatment that includes the primary aging of theturbine engine component at 1975-2050° F. for 2-4 hours and secondaryaging at 1300-1500° F. for 16-24 hours.

To avoid a recrystallization of the turbine engine componentsmanufactured by the hot forming, the service temperature of theseturbine engine components is selected below of the temperature of theprimary aging.

In accordance with the an embodiment, a method of manufacturing ofturbine engine components comprising the step of a fusion weldingpreferably selected from among a laser beam, plasma arc, micro plasma,and electron beam and gas tungsten arc welding, by a melting anddeposition of a powder mix comprising at least two dissimilar nickel andcobalt based powders in quantities of (70-80) wt. % and (20-30) wt. %respectively in a welding pool, wherein the nickel based powdercomprises by wt. %:

Chromium from 6 to 8%,

Cobalt from 6 to 12%,

Molybdenum 1.3 to 1.6%,

Tungsten from 4.5 to 5%,

Tantalum from 2.0 to 6.000,

Niobium from 1 to 4.0%,

Tantalum plus Niobium from 3.0 to 7.0%,

Aluminum from 3.0 to 6.5%,

Hafnium from 0.2 to 1.5%,

Rhenium from 2.5 to 3%,

Germanium from 0 to 1.0%,

Silicon from 0 to 1%,

Yttrium for 0 to 0.2%,

Boron from 0 to 0.015%,

Carbon from 0.01 to 0.1%, and

Ni with impurities to balance, and

And the cobalt based powder comprises by wt. %

Nickel from 10 to 18%,

Chromium from 19 to 21%,

Tungsten from 8 to 10%,

Aluminum from 3 to 6.5%,

Germanium from 0 to 1.000,

Silicon from 0 to 1%,

Yttrium from 0 to 0.450%,

Hafnium from 0 to 1.500

Niobium from 0 to 4%,

Carbon from 0.01 to 0.2% and

Co with impurities to balance;

By progressively moving and solidifying of the welding pool as per apreprogrammed welding path, thereby forming a welding bead with thechemical composing same as the superalloy of the present invention; postweld heat treatment selected from among the high isostatic pressure,annealing, aging or combination of the annealing and aging; machining toa required geometry, and non-destructive testing.

To execute an embodiment above, the powder mix is selected from among apre-alloyed powder blend comprising the dissimilar nickel and cobaltbased powders or nickel and cobalt based powders that are mixed in thewelding pool directly during welding.

BRIEF DESCRIPTION OF THE DRAWINGS

The patent or application file contains at least one drawing executed incolor. Copies of this patent or patent application publication withcolor drawing(s) will be provided by the Office upon request and paymentof the necessary fee.

Reference will now be made, by way of example, to the accompanyingdrawings which show example embodiments of the present application, andin which:

FIG. 1 is the microstructure of the invented cast superalloy in annealedand aged condition depicting: (a) Formation of zigzagged grainsboundaries during solidification; (b) Precipitation of the cuboidal γ′phase during the aging heat treatment;

FIG. 2 is the microstructure of the extruded rods in the aged conditiondepicting: (a) Formation of equiaxed grains with straight boundariesduring extrusion and primary recrystallization; (b) Precipitation of theγ′ phase during the aging heat treatment;

FIG. 3 is the microstructure of LBW welds produced at a room temperaturedepicting: (a) Formation of micro cracks in René 142 weld produced usingGTAW welding with preheating to 1700-1800° F.; (b) The defect freemultilayer weld buildup produced at an ambient temperature using LBWwith the welding powder manufactured from the invented superalloy;

FIG. 4 is the microstructure of the defect free multilayer weld buildupproduced using the LBW at an ambient temperature on the PWA1484 SXsubstrate (base metal) wherein: (a) The crack free fusion of weld andbase metals in as welded condition; (b) Precipitation of the γ′ phase inthe weld metal after the PWHT aging heat treatment;

FIG. 5 is the fracture and EDS mapping (distribution) of some alloyingelements in the tensile sample manufactured from the weld metaldepicting interdendritic precipitation of fine cuboidal Ta—Hf basedintermetallic particles: (a) Ductile fracture of the weld metal tensiletest sample produced using SEM; (b) Distribution of tantalum; (c)Distribution of hafnium;

FIG. 6 is the fractography tensile test sample manufactured from thegermanium free embodiment of the invented superalloy wherein: (a)Fractograph depicting the ductile dimple fracture of the tensile sampleand cuboidal Ta—Hf based intermetallic particles at the bottom ofdimples; (b) The same as a) with higher magnification depictingselecting and marking of typical particles (Spectrum 1 and 2) for EDS;(c) Chemical analysis of the particle marked Spectrum 1 and chemicalcomposition of the selected particle comprising 46.5% Ta-37.3% Hf-9.5%Ni-4.1% Co-1.8% Cr;

FIG. 7 is the microstructure of the weld produced using the inventedsuperalloy on the René 80 substrate wherein: (a) Dendritic structureformed in the weld in ‘as welded’ condition; (b) Microstructure of theweld metal and base material adjacent to the fusion line after theannealing and aging PWHT as per the preferable embodiment;

FIG. 8 is the fractograph of the weld metal test sample subjected to thebend test at an ambient temperature depicting the ductile fracture ofthe sample;

FIG. 9 is the fractograph of the weld sample manufactured from theinvented embodiment of the invented superalloy comprised of 0.85 wt. %germanium and subjected to the tensile testing at 1800° F. depicting:(a) Alternation of a morphology of Ta—Hf based intermetallic particles;(b) Same as a) at higher magnification with the selection of typicalTa—Hf particles for EDS; (c) Mapping of Ta and Hf on the surface of theparticle marked Map Data 19 in FIG. 9a depicting significant enrichmentof this particle with Ta and Hf;

FIG. 10 is the microstructure of LBW weld produced using the powderblend comprising dissimilar nickel and cobalt based powders depicting:(a) Formation of dendritic structure during solidification of a weldingpool; (b) Dissolution of dendrites during the homogenizing annealingfollowed by the aging as per the preferable embodiment;

FIG. 11 is the weld sample manufactured from the alloy 4285 described inthe Example 6 utilizing LBW and 3D AM concept depicting: (a) Manual buttGTAW weld of test samples produced by LBW; (b) Epitaxial dendrite growthin multi pass LBW welds and typical thickness of the LBW weld depositsof 795 μm per pass; (c) Precipitation of cuboidal gamma prime and doublegamma prime phases during post weld heat treatment of the GTAW weldjoint.

Similar reference numerals may have been used in different figures todenote similar components.

Standard Acronyms and Major Definitions

-   -   ASTM—American Society for Testing and Materials (standards)    -   HPT—High Pressure Turbine    -   LPT—Low Pressure Turbine    -   NDT—Non-Destructive Testing    -   NGV—Nozzle Gide Vane    -   PWHT—Post Weld Heat Treatment    -   UTS—Ultimate Tensile Strength    -   SRT—Stress Rupture Test    -   LBW—Laser Beam Welding    -   MPW—Micro-Plasma Welding    -   GTAW—Gas Tungsten Arc Welding    -   EBW—Electron Beam Welding    -   PAW—Plasma Arc Welding    -   SX—Single Crystal Material    -   BM—Base Material    -   3D AM—Three-Dimensional Additive Manufacturing    -   SEM—Scan Electron Microscope    -   EDS—Energy-Dispersive X-ray Spectroscopy    -   IPM—Inch per Minute    -   FPI—Fluorescent Penetrant Inspection

Nickel Based Superalloys—are metallic materials that are used for amanufacturing of turbine engine components and other articles thatexhibit excellent mechanical strength and resistance to creep (tendencyof solid materials to slowly move or deform under stress) at hightemperatures, up to 0.9 melting temperature; good surface stability,oxidation and corrosion resistance. Precipitation strengtheningsuperalloys typically have a matrix with an austenitic face-centeredcubic crystal lattice with precipitation of nickel-aluminum ortitanium-aluminum based γ′ phase. Superalloys are used mostly formanufacturing of turbine engine components.

Hot Forming—Hot forming, which is also known as a hot working, is aprocess in which a metal is shaped under pressure at a fairly hightemperature at which material has sufficient ductility.

High Gamma Prime Nickel Based Superalloys—are nickel-based superalloyscomprising from 3 wt. % to 12 wt. % either aluminum or titanium or totalaluminum and titanium alloying elements.

Laser Beam (Electron Beam, Gas Tungsten Arc, and Plasma Arc) Welding—isa welding process that produces coalescence of materials with the heatobtained from the application of concentrated coherent light beam(electron beam or electric arc respectively) impinging upon the joint orbase material with or without welding material.

Weldability—ability of a material to be welded under imposed conditionsinto a specific, suitable structure and to perform satisfactorily forits intended use.

Structural Turbine Engine Components—various cases, frames, nozzle guidevane rings and other stator parts that ensure engine integrity inservice conditions.

Base Material—is the material of the engine components and test samples.

Energy-dispersive X-ray spectroscopy (EDS)—is an analytical techniqueused for the elemental analysis or chemical characterization of asample.

DESCRIPTION OF EXAMPLE EMBODIMENTS

The invented material belongs to the precipitation strengthening high γ′superalloys and comprises high amount of aluminum, which is the majorwell-known gamma prime forming elements.

The unique combination of strength, ductility, oxidation resistance andweldability is attributed to a precipitation of large volume of highstrength intermetallic gamma prime (γ′) Ni₃Al, double gamma prime (γ″)Ni₃Nb phases, and Ta—Hf—W—Si cuboidal intermetallic particles shown inFIGS. 1b , 9, 10, and 11 in the austenitic ductile γ phase matrix, whichis a solid solution of Co, Cr, Mo, W, Re in nickel, with optimized ratioof all alloying elements. It was found that the fraction volume of γ′and γ″ phases in the invented superalloy exceeding 50 vol. % in the agedcondition.

Ingots for the evaluation of mechanical properties of the inventedsuperalloy were produced by a triple arc re-melt in argon followed bythe annealing and aging heat treatment as per a particular embodiment.

Welding wire was manufactured by the multi-step extrusion of ingots attemperatures 1600-1800° F. followed by pickling for removing of surfaceoxidation.

Welding powder of 45 μm in diameter was produced by gas atomizing ofingots in argon.

To help maximize mechanical properties of the invented precipitationstrengthening superalloy, heat treatment that includes thehomogenization annealing within a temperature range from 2190° F. to2290° F. for 1-2 hours, followed by the primary aging within atemperature range from 1975° F. to 2050° F. for 2-4 hours and thesecondary aging within a temperature range from 1300° F. to 1500° F. for16-24 hours, was developed. This heat treatment was different from theheat treatment frequently used for the heat treatment of R142superalloy, refer to W. Ross and Kevin S. O'Hara for René 142 in “Rene142: High Strength, Oxidation Resistance DS Turbine Airfoil Alloy”,Superalloys 1992, pp. 257-265.

Parameters for PWHT heat treatment of turbine engine components dependon applications. In one embodiment, it was found that the heat treatmentparameters for HPT, LPT NGV and other non-rotating components of turbineengines manufactured by casting and 3D AM comprises annealing within thetemperature range from 2250-2290° F. for 2 hours followed by the primaryaging at 1100-1120° F. for 2 hours and the secondary aging at atemperature of 1480-1500° F. for 24 hours.

PWHT parameters for the heat treatments of HPT and LPT turbine bladesmanufactured from single crystal superalloys and/or repaired by weldingusing the invented welding wire or welding powder includes primary andsecondary aging with the temperature range from 1975° F. to 1995° F. for4 hours and 1300° F. to 1325° F. for 16 hours respectively to avoidrecrystallization of the base material. The heat treatment of turbineengine components manufactured from the invented superalloy by the hotforming comprises also only the primary and secondary aging using theabove disclosed parameters to prevent recrystallization of the basematerial.

Service temperature of the turbine engine components manufactured fromthe invented superalloy by the hot forming was selected below theprimary aging temperature, aiming to exclude recrystallization anddegradation of mechanical properties of the base material in serviceconditions.

Annealing of ingots prior to extrusion or after manufacturing of turbineengine components by casting as per the preferable embodiment results inthe homogenization while aging plays the key role in the formation ofsuperior strength due to a precipitation of γ′ phase. Further,preferable embodiments are explained in more details by examples.

Example 1

To demonstrate the unique combination of high strength and ductility ofthe developed superalloy, samples manufactured from René 142 (R142) andMerl 72 (M72), invented superalloy with the preferable embodiments(samples marked 4275A, 4275B, 4275C, and 4275D), and superalloy with thechemical composition deviated from the preferable embodiment (samplemarked 427X) shown in Table 1, were produced by the triple arc re-meltin argon followed by the homogenization annealing at 2215-2230° F. for 2hours, primary aging at 2035-2050° F. for 2 hours, and secondary agingat 1155-1170° F. for 24 hours.

Test specimens of 0.255-0.275 inch in diameter were machined from ingotsand subjected to the radiographic examination as per ASTM E192-04.Linear indications and pores exceeding 0.002 inch in size were notpermitted. Subsized test samples with the gauge diameter of 0.176-0.180inch and 1.8 inch in length were machined as per ASTM E-8. Tensile testswere conducted as per ASTM E-21 at the temperature up to 1800° F.

Table 1 Chemical Composition of Nickel Based Superalloys with Ni toBalance Samples Ni Cr Co Ta Al W Mo Re Hf C B Y Ge Si Nb M72 15 20 Bal 34.4 9 — — 1 0.35 — 0.45 — — — R142 Bal 6.8 12 6.3 6.1 4.9 1.5 2.8 1.20.12 0.015 — — — — 4275A Bal 9 20 6.0 5.5 5.5 1.0 1.5 0.2 0.10 0.01 0.15— 0.01 — 4275B Bal 10 21.5 5.4 6.0 5.0 1.2 2.5 1.2 0.12 0.01 — — 0.12 —4275C Bal 9.8 20.4 5.4 5.5 5.1 1.2 2.3 1.1 0.14 0.015 0.01 0.85 — —4275D Bal 10.2 22 2.0 4.2 5.5 1.2 3.5 1.5 0.12 0.01 0.1 0.2 0.8 — 4275EBal 10.1 22 5.45 5.7 5.95 2 2.1 1.15 0.13 0.01 0.11 — 0.1 — 427X Bal 1026 5.5 6.2 5.4 1.4 2.0 1.1 0.12 0.01 0.1 — — — 4287 Bal 10 18 2 6.2 5.61.2 2.2 0.3 0.1 0.015 — — 0.5 4

Solidification of ingots resulted in the formation of zigzagged grainsboundaries shown in FIG. 1a , which enhances mechanical properties ofthe developed superalloy. The post weld (PWHT) aging heat treatmentresults in the precipitation of high volume of γ′ phase shown in FIG. 1b.

Precipitation of a large volume of high strength γ′ phase in the ductileaustenitic matrix results in a formation of the desirable combination ofhigh strength and ductility as shown in Table 2. Ductility (elongation)of the invented superalloy is superior to the ductility of standard R142samples while strength is superior to M72.

TABLE 2 Mechanical Properties of Ingots Produced by Arc Triple Re-Meltin Argon Test Temp. UTS, 0.2% Yield Elong. Material ° F. KSI Strength,KSI % M72 1800 23.1 15.7 86.8 R142 1800 71.2 70.5 1.0 4275A 70 172.1142.0 7.0 4275A 1450 136.7 125.8 8.6 4275A 1600 113.3 93.1 6.9 4275A1800 70.9 61.7 9.8 4275B 1800 71.5 68.5 5.0 4275D 1800 63.6 55.0 14.0427X 1800 43.7 37.8 18.2

Example 2

Low γ′ wrought AMS 5664 Inconel 718 (IN718) and AMS 5704 Waspaloysuperalloys have been used for the manufacturing of structural turbineengine components due to high strength at the temperature up to 1200° F.and good workability. However, further heating of IN718 and Waspaloy to1800° F. drastically reduced strength and stress rupture properties(SRT) of these superalloys as shown in Table 3.

Due to a good combination of strength at the temperature up to 1800° F.and workability of the developed high gamma prime superalloys, it isfound that the developed high gamma prime superalloys are most prominentfor a substitution of standard wrought superalloys for a manufacturingof structural turbine engine components utilizing hot forming processes.To evaluate mechanical properties of the invented superalloy in wrought(hot formed) condition, ingots were subjected to the extrusion as perthe preferable embodiment to produce bars of 0.225 inch in diameter,which further were subjected to the primary aging at the temperature of1950° F. for 4 hours and secondary aging at 1300° F. for 24 hours.

The subsized test samples of 1.8 inch in length with the gauge diameterof 0.158-0.162 inch were machined as per ASTM E-8. Tensile tests wereconducted as per ASTM E-8 at 70° F., and as per ASTM E-21 at 1200° F.and 1800° F. The stress rupture testing was conducted at temperatures of1200° F., 1350° F., and 1800° F. as per ASTM E-139.

Extrusion of the invented superalloy at high temperature resulted in aformation of the equiaxed structure with the straight grain boundariesshown in FIG. 2a , which were different from zigzagged boundaries formedduring the solidification of ingots shown in FIG. 1a . The primary agingheat treatment resulted in a precipitation of γ′ phase shown in FIG. 2b.

As it was found by experiments, UTS and SRT properties of the developedsuperalloy were superior to UTS and SRT of Inconel 718 and Waspaloy upto 1800° F. as shown in Table 3 and 4 respectively.

TABLE 3 Tensile Properties of Wrought (Hot Formed by Extrusion)Superalloys Test Temp. UTS, 0.2% Yield Elongation. Material ° F. KSIStrength, KSI % Inconel 718 70 186.3 161.2 12.5 1200 162.5 138 10.5 180015.7 8.5 67.9 Waspaloy 70 195.7 168.3 16 1200 186.4 139.5 20.4 1800 30.121.5 49.9 4275A 70 182.5 155.6 10.5 1200 174.2 145.7 11.0 1800 59.6 43.35.1

TABLE 4 SRT Properties of Hot Formed (Extruded) Rods Test Temp.Stresses, Time to Material ° F. KSI Rupture, Hours Inconel 718 1200 10028 1800 15 1.4 Waspaloy 1350 80 26.5 1800 15 4.3 4275A 1200 100 232 135080 447.8 1800 15 31.2

Combination of high strength, ductility and workability makes theinvented superalloy most prominent for a manufacturing of turbine enginecomponents by the hot forming.

Example 3

To simulate the repair of turbine engine components manufactured fromsingle crystal materials using manual GTAW and automatic LBW welding,test samples were produced using the developed superalloy in a form ofwelding wire and welding powder respectively, and using standard René142 welding wire for GTAW with preheating to 1700-1800° F. and LBW at anambient temperature.

Preheating was used for GTAW with René 142 welding wire to producesamples for tensile and SRT testing because welding at an ambienttemperature results in extensive cracking of René 142 welds as shown inFIG. 3 a.

Multi pass LBW with welding powder manufactured from the inventedsuperalloys and GTAW with welding wire manufactured from the inventedsuperalloys were performed at an ambient temperature so as to produceweld samples marked LBW4275 and GTAW4275. Welds were free of cracks.Typical microstructure of these samples is shown in FIG. 3b and FIG. 4a.

The post weld heat treatment of welds included the homogenizationannealing at 2200° F. for two hours followed by the primary aging at1975-1995° F. for 4 hours and the secondary aging at 1300-1320° F. for16 hours to exclude recrystallization of HPT blades manufactured fromthe PWA1484 SX material, which resulted in a precipitation of γ′ phaseshown in FIG. 4b with the fraction volume of 49.2 vol. %.

Flat subsized ‘All Weld Metal’ samples of 0.050 inch in thickness wereproduced as per ASTM E-8 and subjected to the tensile testing at 1800°F. as ASTM E-21 and SRT at 1800° F. and stresses of 22 KSI as per ASTME-139.

TABLE 5 Tensile and Creep Properties of Rene 142 and 4275 Weld MetalsWeld Method Test 0.2% Yield Time to and Sample Temp. UTS, Strength,Elong. Rupture ID ° F. KSI KSI % in Hours GTAW R142 1800 34.8 34.0 2.724.2 LBW4275B 1800 71.7 52.6 6.5 278.5 GTAW4275B 1800 67.5 53.8 8.7216.8

As follows from Table 5, ductility and SRT properties of LBW and GTAWwelds produced from the invented superalloy were superior to propertiesof standard René 142 welds.

Low tensile and SRT properties of René 142 welds were attributed to aformation of microcracks shown in FIG. 3 a.

High tensile and creep properties, as well as good ductility andweldability of the developed superalloy, were attributed to theprecipitation of high volume of high strength cuboidal γ′ phase in theductile Ni—Cr—Co—Re—W—Mo solid solution of gamma matrix andinterdendritic precipitation of fine cuboidal Ta—Hf based intermetallicparticles shown in FIGS. 5 and 6.

Example 4

Germanium has not been used for a manufacturing of Ni based superalloysdespite that nickel based brazing material comprising Ni-(5-40) wt. %Cr-(15-40) wt. % Ge as per the U.S. Pat. No. 2,901,374 was invented in1954. Despite that germanium is the melting point depressant that shouldaffect high temperature strength, we discovered that the addition of upto 0.85 wt. % of germanium to the invented superalloys, which was marked4275C in Table 1, improves weldability and produced defect free welds onthe René 80 as shown in FIG. 7.

Welding of test samples was done manually with the weld current of 75-80A, voltage of 9-10 V and welding speed of 1-1.2 ipm (inch per min).After welding, samples were subjected to heat treatment that includedannealing at 2190° F. for 2 hours, primary aging at 1975° F. for 2 hoursfollowed by the secondary aging at 1550° F. for 16 hours. The tensilesamples for testing were machined as per ASTM E-8 from the base materialand weld, and subjected to tensile testing at 1800° F.

The weld metal was also subjected to the semi guided bend test as perASTM E-190 at an ambient temperature.

In addition to above, the cylindrical samples manufactured from the René80 and invented superalloy were subjected to the cyclic oxidationtesting at 2050° F. in 500 hours. Duration of each cycle was 1 hour thatincluded exposure to 2050° F. for 50 min followed by cooling to about700° F. and reheating to 2050° F. for 10 min.

As it was found by experiments, the strength and oxidation resistancesof welded joints and weld metal were superior to the René 80 basematerial as shown in Tables 6A and 6B.

TABLE 6A Tensile Properties of the René 80 and Invented Superalloy Test0.2% Yield Weld Method, and Temp. UTS, Strength, Elong. Material ° F.KSI KSI % René 80 1800 55.3 45.3 16.5 René 80—4275C 1800 61.8 48.1 12.2Dissimilar Welded Joint

TABLE 6B Oxidation Properties of the René 80 and. Invented superalloy at2050° F Weld Method, Weight Lost in gram after and Material exposure inair for 200 hours René 80 3.1583 4275C Weld Metal 0.0028

Bend samples produced from the weld metal fractured approximately at90°, demonstrating unique ductile of the invented superalloy as shown inFIG. 8 that was not reported on any welds produced on known high γ′superalloys. As it was found by experiments, germanium enhances bondingbetween Ta—Hf intermetallic particles and changes morphology of theseparticles as shown in FIG. 6a and FIG. 9a respectively. The EDS analysisconfirmed that particles were produced by Ta—Hf based intermetalliccompound, refer to FIGS. 9b and 9c . This effect was unknown because onthe contrary to Si, which belongs to the same IVA group of chemicalelements, germanium within the specified range does not result in theformation of the intergranular and interdendritic Ni—Ge based eutecticsthat affect mechanical properties of Si bearing nickel-basedsuperalloys.

Therefore, superior mechanical properties of the Ge-bearing embodimentof the invented superalloy were achieved by the combination of highcontent of γ′ phase, and strengthening of grain and dendrites boundariesby fine Ta—Hf based intermetallic particles with coherent bonding withthe ductile Ni—Cr—Co—Re—Mo—W based matrix shown in FIG. 9a , andpeculiarities of a solidification of a welding pool, which is producedby the dissimilar nickel and cobalt based powders that are meltedtogether in the welding pool and then solidified, produces properties ofwelds superior to properties of welds produced by using homogeneouswelding powders and wires. Oxidation resistance was enhanced by theoptimized content of Cr, Al, Si in a combination with Ge and all otheralloying elements of the invented superalloy.

Based on the test results, the welding wire and powders manufacturedfrom the invented superalloy were found most prominent for the tiprepair of HPT and LPT blades, ensuring the optimal clearance between thetip of blades and stator, low fuel consumption, and high efficiency ofturbine engines through the full engine cycle between overhauls.

Example 5

To demonstrate 3D AM process for a manufacturing of turbine enginecomponents, samples of 4 inch in length by 1 inch in height and 0.125inch in thickness were produced, using the LAWS1000 laser welding systemequipped with 1 kW IPG laser and two powder feeders allowing mixing oftwo dissimilar nickel and cobalt based dissimilar powders directly inthe welding pool as well as performing welding using the pre-alloyedpowder blend.

The example below is depicting welding with the pre-alloyed powder blendthat comprises 75 wt. % of the nickel based powder and 25 wt. % of thecobalt based powder. The nickel-based powder comprises 6.8 wt. % Cr, 12wt. % Co, 1.5 wt. % Mo, 4.9 wt. % W, 6.3 wt. % Ta, 6.1 wt. % Al, 1.2 wt.% Hf, 2.8 wt. % Re, 0.1 wt. % Si, 0.12 wt. % C, 0.01 wt. % B, and Ni tobalance. The cobalt based powder comprises 17 wt. % Ni, 20 wt. % Cr, 3wt. % Ta, 9 wt. % W, 4.4 wt. % Al, 0.45 wt. % Y, 0.1 wt. % Si, 0.015 wt.% C and Co to balance.

Welding parameters that were used to produce samples are provided below:

Laser beam power—480 W (Watt)

Deposition rate—3.8 g/min (gram per min)

Welding speed—3.5 ipm (inch per min)

Beam oscillation speed across the weld—40 imp

Inert gas—argon

During multi pass weld deposition, the welding pool was movedprogressively as per the preprogrammed welding path with the speed of3.5 ipm, which, due to solidification, results in the formation of awelding bead with the preferable chemical composition that is same asthat of the invented superalloy. Chemical composition of the weld metalsample marked 4275E is provided in Table 1.

After welding test, samples were subjected to the primary aging at2035-2050° F. for 2 hours, and secondary aging at 1155-1170° F. for 24hours, machining to a required geometry followed by a non-destructivetesting that includes FPI as per AMS 2647 and radiographic inspection asper ASTM E192-04. Weld discontinuities that exceeds 0.002 inch in sizewere not permitted.

Subsized test samples were produced from welds as per ASTM E-8 andsubjected to tensile testing at 1775° F. as per ASTM E-21.

Welding resulted in a formation of dendritic structure with theepitaxial grain growth as shown in FIG. 10a . Welds were free of cracksand other weld discontinuities.

The post weld homogenizing and aging heat treatment resulted inprecipitation of large volume of gamma phase as shown in FIG. 10b .

TABLE 7 Tensile and SRT Properties of Welds Produced by LBW with thePowder Blend Weld Metal Test UTS, 0.2% Yield Elong. Sample ID Temp. ° F.KSI Strength, KSI % 4275E 1775° F. 74.8 63.5 7.4

As follows from the Table 7, weld samples demonstrate excellent strengthand good ductility at a temperature of 1775° F., despite the balkcontent of Al in weld metal of 5.7 wt. %.

Superior weldability, strength and ductility of the invented superalloythat comprises 5.7 wt. % of aluminum were achieved by the peculiaritiesof a solidification of the welding pool produced by the dissimilarnickel and cobalt based powders.

Known nickel-based superalloys comprising 5.7 wt. % Al are not weldableat an ambient temperature, while LBW welding using the mix of dissimilarpowders and/or powder blends, which due to a solidification of a weldingpool forms welds with the balk chemical composition corresponding to thechemical composition of the invented superalloy, produces sound weldswith high mechanical properties.

Example 6

To simulate the repair of turbine engine components manufactured fromRene N5 single crystal (SX) material, LBW welding at an ambienttemperature was performed on the SX substrate using the powder blend of78-80% Ni-based powder ‘A’ and 20-22% of Co-based powder ‘B’ comprising:

Powder A: Ni-8% Cr-8% Co-1.5% Mo-4.5% W-3.5% Ta-6% A-0.75% Hf-0.15%Si-3.5% Re-1.2% Nb-0.012% B-0.1% C

Powder B: Co-18% Cr-15% Ni-10% W-0.05% Hf-3.5% A-0.15% Si-0.1% C

Dissimilar Ni based powder A and cobalt based powder B were melted bythe laser beam in a welding pool producing a homogeneous alloy. Due to asolidification of a welding pool the welding bead with the bulk chemicalcomposition below (material 4285) was formed:

Material 4285: Ni-10% Cr-16% Co-1.2% Mo-5.6% W-2.8% Ta-1% Nb-5.5%Al-0.5% Hf-0.15% Si-0.01% B-0.1% C-2.2% Re

After welding, test samples produced by LBW utilizing 3D AM concept weresubjected to the primary aging at 1975° F. for 4 hours followed thesecondary aging at 1650° F. for 4 hours.

After welding and heat treatment, samples were subjected tometallographic and radiographic examinations. No cracks and other welddiscontinuities were found.

Subsized tensile specimens were machined from the weld samples andsubjected to the tensile testing at an ambient temperature todemonstrate high ductility of the developed material as per ASTM E-8alloy. Tensile properties of 4285 all weld metal samples are providedbelow:

-   -   UTS=187.2 KSI    -   0.2% Yield Strength=173.4 KSI    -   Elongation=11.2%

Additional test was carried out on the evaluation of weldability of 4285material

As it was found by experiments, GTAW welds produced using manual weldingat an ambient temperature on 4285 material produced by automatic LBWwere free of weld discontinuities as shown in FIG. 11 a.

Multi pass LBW welding resulted in a formation of the dendriticstructure and epitaxial grain growth as shown in FIG. 11 b.

Post weld heat treatment resulted in a precipitation of the formation offine γ′ Ni₃Al, double gamma prime γ″ Ni₃Nb and intermetallic Ta—Hf—W—Sistrengthening phases in a ductile austenitic matrix of the GTAW weldmetal as shown in FIG. 11c enhancing ductility, weldability and highstrength of the invented superalloy.

GTAW butt welded joints shown in FIG. 11a at demonstrated high strengthat 1400° F.:

-   -   UTS=142.1 KSI    -   0.2% Yield Strength=132.9 KSI    -   Elongation=5.2%

Good weldability, high ductility, and enhanced strength of the 4285superalloy were attributed to the optimized with the range from 3 to 7wt. % of the total content of Ta and Nb in the invented superalloy.

Example 7

To simulate the repair of turbine engine components manufactured fromRene N5 SX material, LBW welding at an ambient temperature was performedon the SX substrate using the powder blend of 70-72% of Ni-based powder‘C’ and 28-30% of Co-based powder ‘D’ comprising:

Powder C: Ni-6% Cr-6% Co-1.7% Mo-5.6% W-3.4% Ta-4% Nb-6.2% A-0.3%Hf-0.5% Si-3% Re-0.02% B-0.1% C

Powder D: Co-21% Cr-5.6% W-4% Nb-6.2% A-0.3% Hf-0.5% Si-0.1% C.

The welding bead with the bulk chemical composition of the material 4287below was formed due to a solidification of a welding pool:

Material 4287 comprises:

Ni-10% Cr-18% Co-1.2% Mo-5.6% W-2% Ta-4% Nb-6.2% A-0.3% Hf-0.5%Si-0.015% B-0.1% C-2.2% Re

The post weld heat treatment of welds included the primary aging at1975° F. for 4 hours followed by the secondary aging at 1650° F. for 4hours.

After welding and heat treatment samples were subjected tometallographic and radiographic examinations. No cracks and other welddiscontinuities were found.

Subsized tensile specimens were machined from the weld samples as perASTM E-8 and subjected to the tensile testing at 1800° F. as per ASTME-21. Tensile properties of 4287 all weld metal samples are providedbelow:

-   -   UTS=63.6 KSI    -   0.2% Yield Strength=49.3 KSI    -   Elongation=14.5%

While the invention has been described in terms of preferableembodiments, it is apparent that other forms of the current inventioncould be adopted by one skilled in the art. Therefore, the scope of theinvention is to be limited only by the claims herein.

Certain adaptations and modifications of the described embodiments canbe made. Therefore, the above discussed embodiments are considered to beillustrative and not restrictive.

What is claimed is:
 1. A gamma prime nickel-based superalloy, comprisingby wt. %: Chromium from 9.0 to 10.5%, Cobalt from about 16 to 22%,Molybdenum from 1.0 to 1.4%, Tungsten from 5.0 to 5.8%, Tantalum from2.0 to 6.0%, Niobium from 2.2 to 4.0%, Tantalum plus Niobium from 4.2 to7.0%, Aluminum from 3.0 to 6.5%, Hafnium from 0.2 to 1.5%, Germaniumfrom 0 to 1.0%, Yttrium from 0 to 0.2%, Silicon from 0.2 to 1.0%, Boronfrom 0 to 0.015%, Carbon from 0.01 to 0.2%, Rhenium from 1.5 to 3.5%,and Nickel with impurities to balance.
 2. The gamma prime nickel-basedsuperalloy according to claim 1 wherein the total content of germaniumand silicon is within 0.9-1.1 wt. %.
 3. A use of the gamma primenickel-based superalloy according to claim 1 as the material for awelding wire, welding powder, or turbine engine components.
 4. A gammaprime nickel-based superalloy as per claim 1, comprising by wt. %:Chromium from 9.5 to 10.5%, Cobalt from about 17 to 19%, Molybdenum from1.1 to 1.3%, Tungsten from 5.0 to 5.8%, Tantalum from 1.5 to 2.5%,Niobium from 2.5 to 4.0%, Tantalum plus Niobium from 5.0 to 6.0%,Aluminum from 6.0 to 6.5%, Hafnium from 0.3 to 0.5%, Silicon from 0.4 to0.8%, Boron from 0.01 to 0.015%, Carbon from 0.01 to 0.12%, Rhenium from2.0 to 2.5%, and Nickel with impurities to balance.